Boosting the stability of perovskites with exsolved nanoparticles by B-site supplement mechanism

Perovskites with exsolved nanoparticles (P-eNs) have immense potentials for carbon dioxide (CO2) reduction in solid oxide electrolysis cell. Despite the recent achievements in promoting the B-site cation exsolution for enhanced catalytic activities, the unsatisfactory stability of P-eNs at high voltages greatly impedes their practical applications and this issue has not been elucidated. In this study, we reveal that the formation of B-site vacancies in perovskite scaffold is the major contributor to the degradation of P-eNs; we then address this issue by fine-regulating the B-site supplement of the reduced Sr2Fe1.3Ni0.2Mo0.5O6-δ using foreign Fe sources, achieving a robust perovskite scaffold and prolonged stability performance. Furthermore, the degradation mechanism from the perspective of structure stability of perovskite has also been proposed to understand the origins of performance deterioration. The B-site supplement endows P-eNs with the capability to become appealing electrocatalysts for CO2 reduction and more broadly, for other energy storage and conversion systems.

nanoparticle exsolution of B site cations in perovskite lattice leaves many B site vacancies and causes detrimental A site segregation, leading to the losses of B(n+1)-O-Bn+ pathway. In this work, authors have insisted that the guest Fe metal coated on Sr2Fe1.3Ni0.2Mo0.5O6-δ (SFNM) would fill the B site vacancies by topotactic ion exchange (TIE) mechanism and mitigate the Sr segregation on the surface, thus enhancing structure stability and facilitating stable CO2 electrolysis performance at high voltage. Several characterizations such as XRD, SEM, TEM, and elemental mappings were performed to verify the role of guest Fe, filling the B site vacancies during exsolution process and suppressing Sr segregation on perovskite surface. In potential step chronoamperometry, SFNM without guest Fe (SFNM+0.0Fe-red) showed significant decay of current density (j), whereas SFNM with guest Fe metals (SFNM+1.2Fe-red) showed stable j profiles, especially under high operating voltage (1.8V).
In this work, the results are noteworthy and would be useful to develop a highly active CO2 electrolysis cathode catalyst with superior stability in SOEC system. Authors have provided detailed mechanisms of perovskite structure evolution during the exsolution and TIE-assisted exsolution, and most of conclusions are supported by the results. It might be suitable to be published in this journal after addressing below issues and comments: 1. In page 3~4, authors have mentioned that the exsolution process on the perovskite leaves many B site vacancies within perovskite bulk and thus A site segregation on the surface as below equation:

A_(1-α)BO_(3-δ) → (1-α)ABO_(3-δ')+αB
A-site deficient design has been adopted in numerous other researches of SOEC CO2 cathodes, especially in [D. Neagu et al., Nat. Chem., 2013, 5, 11, 916-923]. It might be helpful to mention on your manuscript that what is the strong point of your strategy to coat additional guest Fe metals on perovskite surface to induce TIE-assisted exsolution compared to the one, which controls the A site nonstoichiometry of perovskite oxide. elevated temperatures 5 . Therefore, the perovskites with exsolved nanoparticles (P-eNs) provide an appealing platform for the large-scale energy storage and conversion technology compared to their nanoparticle-free counterparts.
It is well known that many factors such as the types of cathode, electrolyte and anode could contribute to the higher current density of Sr2Fe1.4Mn0.1Mo0.5O6 perovskite based SOEC in reference [J. Mater. Chem. A, 2019, 7, 11967] than Sr2Fe1.3Ni0.2Mo0.5O6 (SFNM)-based exsolved SOEC in our work. In the reference [J. Mater. Chem. A, 2019, 7, 11967], LSCF-SDC/LSGM/SFMM0.1-SDC was assembled for electrochemical experiments, while LSCF-GDC/GDC/YSZ/GDC/ SFNM+ Fe ( =0.5, 0.8, 1.2)-GDC were fabricated in our study. It is known that the LSGMsupported electrolysis cells perform better than the YSZ-supported ones under the same condition. Therefore, the direct comparison of the cell current density is only meaningful when all the other variables are kept identical. For example, under the same condition, the current density of SFNM+1.2Fe-red-GDC is higher than that of SFNM+0.0Fe-red-GDC, this indicates that the catalytic activity can be enhanced by B-site supplement strategy.
Alternatively, the total polarization resistance (Rp) may be a better indicator for evaluating the catalytic activities of cathode materials since Rp is dominated by the cathode reaction 6 . The Nyquist plots and the fitted RS, Rp (equivalent circuit model LR(QHRH)(QLRL)) of SFNM+1.2Fe-red-GDC at 800 o C are provided in Fig. R1 (also Supplementary Fig. 11) and Table R1 (also Supplementary Table 3). The current density, polarization resistance and CO productivity of SFNM+1.2Fe-red-GDC in our work are compared with those of other state-of-art P-eNsbased SOECs, as shown in Table R2 (also Supplementary Table 4). SFNM+1.2Fe-red-GDC shows the lowest Rp at high voltages (≥ 1.6 V) among the state-of-art P-eNs-based cathode candidates as listed in Table R2, indicative of the higher catalytic activities of SFNM+1.2Fe-red-GDC at higher negative potentials. In addition, the shortterm stability performances of SFNM+0.0Fe-red-GDC and SFNM+1.2Fe-red-GDC are compared with that of other state-of-art P-eNs-based SOECs, as shown in Table R3 (also Supplementary Table 6). SFNM+1.2Fe-red-GDC shows the competitive short-term stability at the voltages ≥ 1.6 V compared with the SFNM+0.0Fe-red-GDC and other state-of-art P-eNs-based SOECs listed in Table R3.
Most importantly, SFNM with varied B-site occupation were selected as the prototypes in our study to highlight the significance of robust perovskite scaffold for enhancing the reactivity and stability of the P-eNs. The faster reaction kinetics and more stable performances of SFNM+1.2Fe-red-GDC in comparison to SFNM+0.0Fe-red-GDC, especially at high voltages, confirm the feasibility of the B-site supplement strategy. It suggests that this strategy would be applied to other P-eNs listed in the Tables R2-R3 to improve their electrochemical performances for CO2 electrocatalysis and more broadly, for other energy storage and conversion systems.    Table R3 (also Supplementary Table 6). Comparison of the short-term stability with the state-of-art P-eNs for pure CO2 electrocatalysis in SOEC.
Note: The values of the current density in the short-term stability of the cited literatures are determined by the Digitizer function of Origin software.
We have revised the manuscript (L4-9, P13 and L1-3, P19) and Supplementary Information (SI) (P35 and P37-38) accordingly, as shown below: Revised manuscript (L4-9, P13): "In addition, the Nyquist plots and the fitted Rp of SFNM+1.2Fe-red-GDC at 800 o C are obtained (Supplementary Fig. 11     We have revised the manuscript (L11-19, P22 and L1-5, P23) and SI (P26-27) accordingly, as shown below: Revised manuscript (L11-19, P22 and L1-5, P23): "To verify the exsolution scenario and the structural decomposition during the electrolysis process, the SFNM without pre-reduction treatment was fabricated as the 9 composite cathode material (SFNM-oxi-GDC) for the electrochemical testing ( Supplementary Fig. 20) and the post-mortem characterizations . As shown in Supplementary Fig. 20d, the SFNMoxi-GDC shows the similar long-term stability profile as the SFNM+0.0Fe-red-GDC. The SEM, SE-STEM and STEM-EDS results confirm that partial Ni elements segregate from the SFNM bulk and the newly born fibrous phases appear on the cathode surface . Meanwhile, the Raman feature peaks of SrCO3, carbon with lattice defect (D band) can be observed ( Supplementary Fig. 23), suggesting that the Ni element insitu exsolves from the SFNM bulk and subsequently results in the carbon fiber growth during CO2 electrolysis. This can be ascribed to the lower co-segregation energy of Ni in SFNM, which causes the structural decomposition under the synergistic effect of 1.6 V and 850 o C." Revised SI (P26-27): "To further clarify the degradation by exsolution, the catalytic activity and stability performances of Sr2Fe1.3Ni0.2Mo0.5O6 (SFNM) based SOEC without exsolution have been evaluated for comparison ( Supplementary Fig. 20). As can be seen from the current density-voltage ( -V) curves at 800 and 850 o C, SFNM-based SOEC without exsolution (SFNM-oxi-GDC) exhibits the lower than those of the SFNM+0.0Fe-red-GDC and SFNM+1.2Fe-red-GDC under the same condition. Additionally, the SFNM-oxi-GDC shows the similar short-term/long-term stability profiles as the SFNM+0.0Fe-red-GDC. In term of the short-term stability, the SFNM-oxi-GDC shows satisfactory stability at 1.0-1.4 V, while the SOEC experiences a rapid degradation when the voltage exceeds 1.6 V. For the long-term stability, the SOEC experiences a significant degradation at the initial stage, followed by a steady -V profile. In fact, the also experiences a slight increase, peaking at 31 h with a maximum of 0.18 A cm -2 . Then, the SFNM-oxi-GDC shows a visible degradation after 60 h electrolysis, and finally destabilized." Comment #2. The HR-TEM results for SFNM+0.0Fe-red and SFNM+0.0Fe-red were clearly showed the difference between the conventional and TIE exsolved nanoparticles. What are the morphologies like for each of them after stability tests?
Authors' response to Comment #2: As suggested by the reviewer, the SEM, SE-STEM and STEM-EDS images are added to characterize the cathode surface morphologies of SFNM+0.0Fe-red-GDC and SFNM+1.2Fe-red-GDC after the long-term stability tests (Figs. R6 and R7 [also Figs. 5d-g in the revised manuscript)]. The obvious morphology changes and coarsening of exsolved particles can be observed on both substrates, and the coarsening of surface nanoparticles on SFNM+1.2Fe-red-GDC is less severe than that on SFNM+0.0Fe-red-GDC (Fig. R6). Furthermore, more pronounced Sr-and Mo-derived phase separation in the form of irregular particles appear on the surfaces of SFNM+0.0Fe-red-GDC, (Fig. R7), suggesting that SFNM+0.0Fe-red-GDC undergoes the more significant phase decomposition than SFNM+1.2Fe-red-GDC under the synergistic effect of the high voltage and CO2/CO mixed atmosphere during the CO2 electrolysis.  We have revised the manuscript (L1-11, P21) accordingly as shown below: Revised manuscript (L1-11, P21): "After the long-term stability tests, the significant coarsening and decrease in the population of nanoparticles can be observed on the cathode surfaces of SFNM+0.0Fe-red-GDC and SFNM+1.2Fe-red-GDC (Figs. 5d-e). Both cathodes were scratched from the cells for the element distribution detection by secondary electron-scanning TEM (SE-STEM) and STEM-EDS. More pronounced Sr-and Mo-derived phase separation in the form of irregular particles appear on the surfaces of SFNM+0.0Fe-red-GDC, (Figs. 5f-g), suggesting that SFNM+0.0Fe-red-GDC undergoes the more significant phase decomposition than SFNM+1.2Fered-GDC under the synergistic effect of the high voltage and CO2/CO mixed atmosphere during the CO2 electrolysis."   Fig. 4a in the revised manuscript)). To simplify the model, we assume that all Ni elements are exsolved from the SFNM substrate after the reduction. Therefore, the binding energy between the Fe-site vacancy and the oxygen vacancy is calculated in the Sr2Fe1.5Mo0.5O6 (SFM). The three defective configurations and corresponding formation energy are provided in Fig. R8 (also Supplementary Fig.   15) and Table R4 (also Supplementary Table 5). The calculated binding energy of ′′ − •• is -1.73 eV indicates that the additional association barrier needs to be overcome to achieve the jumping of oxygen vacancies 15 . In addition, the bulk diffusion constant ( ℎ ) of SFNM+1.2Fe-red obtained by fitting the ECR curve is 2.716 × 10 −5 cm 2 s -1 , much higher than 1.208 × 10 −5 cm 2 s -1 of SFNM+0.0Fe-red 16 . It suggests that there are more available oxygen ion transfer pathways in SFNM+1.2Fe-red, which lead to the higher oxygen ion conductivity and lower charge transport resistance.    Table 5), indicating that the additional association barrier needs to be overcome to achieve the jumping of oxygen vacancies 15 . In contrast, the full occupation of B-sites on SFNM+1.2Fe-red allows the elimination of constraint of B-site defect to surrounding oxygen vacancies. The electrical conductivity relaxation (ECR) experiments for SFNM+0.0Fe-red and SFNM+1.2Fe-red were carried out to study the oxygen ion bulk diffusion 17,18 (Fig. 4a). The bulk diffusion constant ( ℎ ) of SFNM+1.2Fe-red obtained by fitting the ECR curve is 2.716 × 10 −5 cm 2 s -1 , much higher than 1.208 × 10 −5 cm 2 s -1 of the SFNM+0.0Fe-red 16 . It suggests that there are more available oxygen ion transfer pathways in SFNM+1.2Fe-red, which lead to the higher oxygen ion conductivity and lower charge transport resistance." Revised SI (P20-21): "It has been reported that the association (trapping) of oxygen vacancy caused by other defects in perovskite lattice is detrimental to oxygen ion diffusion 15,19,20 . And it is widely accepted that the association of oxygen vacancy can be determined by calculating the binding energy between the two adjacent point defects 19,21 ( ′′ − •• pair in our manuscript). The binding energy ( ) can be calculated by following equation 20 : In our case, To simplify the model, we assume that all Ni elements are exsolved from the SFNM substrate after the reduction. Therefore, the binding energy between the Fe-site vacancy and the oxygen vacancy is calculated in the Sr2Fe1.5Mo0.5O6 (SFM). The three defective configurations and the corresponding formation energy are provided in Supplementary Fig. 15 and Supplementary Table 5. As a result, the calculated binding energy of ′′ − •• is -1.73 eV. It indicates that the existence of B-site vacancies would hinder the transport of oxygen vacancies, thus resulting in the reduction of the free oxygen vacancy population." Comment #4. (minor comment) The caption of Figure 4d and 4e was reversed.

Authors' response to Comment #4:
We thank the reviewer for pointing this out. Authors' response to Comment #5: As suggested by the reviewer, we have moved the following spectra from SI to the revised manuscript including the FTIR spectra of CO2 chemisorption and physisorption for SFNM+0.0Fe-red and SFNM+1.2Fe-red at 600 o C (Fig. 3d), the Raman spectra collected from cathode surface of SFNM+0.0Fe-red-GDC whose stability test was interrupted before entering the Phase II (Fig. 6b). In the revised manuscript, we have also added the normalized electrical conductivity relaxation curves (Fig. 4a), the temperature-dependent electrical conductivities (Fig. 4d), the SEM, SE-STEM and STEM-EDS images of cathode surface microstructure of SFNM+0.0Fe-red-GDC and SFNM+1.2Fe-red-GDC after long-term stability tests (Figs. 5d-5g).

Reviewer #2 (Remarks to the Author):
In this work titled "B-site supplement mechanism: A new approach to boosting stability of perovskites with exsolved nanoparticles", the authors describe the effect of B-site supplement (or topotactic ion exchange) of Fe nitrate guest in double perovskite Sr2Fe1.3Ni0.2Mo0.5O6-δ (SFMN). Unfortunately, I have to say that the electrochemical part of this paper is poor and there are too many flaws in this manuscript. Thus, I can't recommend this manuscript to be published in Nature Communications. Some major comments and questions are listed as below:

Authors' response:
We thank the reviewer very much for carefully reviewing our manuscript and providing insightful comments. To address the concerns raised by the reviewer, we have carefully revised the manuscript according to the reviewer's suggestions, as shown in the following point-to-point responses. We believe that the tendency of our

Authors' response to Comment #1:
It might be possible that the novelty of our work might not have been well perceived and we would like to respectfully reiterate that the significance of our work is to provide the instruction for designing P-eNs with high stability for CO2 reduction and more broadly, for other energy storage and conversion systems. This is of critical importance because the rapid degradation of P-eNs for CO2 electrocatalysis at high voltages greatly impedes their practical applications and this issue has not been well elucidated so far. In this study, we firstly reveal that the structural instability of perovskite scaffold may be the main reason for the rapid degradation of P-eNs for CO2 electrocatalysis at high voltages. Then, the B-site supplement strategy is proposed to yield a robust perovskite scaffold, which leads to higher catalytic activity and stability of the P-eNs at high voltages. Lastly, the detailed degradations of P-eNs with and without B-site supplement are demonstrated. Therefore, the innovation in our work lies in enhancing the stability of P-eNs for CO2 electrocatalysis in SOEC by reinforcing the structural stability of perovskite scaffold rather than the investigation on topotactic ion exchange (TIE) technology, which is a synthetic method for us to verify the feasibility of B-site supplement strategy.
In the past few decades, exsolution has received extensive attentions as an emerging technique for in-situ growing nanoparticles, we have summarized the landmark work of exsolution technology on perovskites in the past ten years 5,22,23 24,25,26,27,28,29,30 in Fig. R10. Most studies have focused on the enhancement of catalytic activity by promoting the exsolution of the nanoparticles and tuning the features of exsolved nanoparticles, while relatively few studies on the stability of P-eNs. Figure R10. The landmark studies of the exsolution on perovskites in the past decade.
Currently, the researches on the stability of P-eNs are mainly focused on the unique socked interface 5,31 . Neagu et al. demonstrated that the exsolved nanoparticles are closely embedded into the parent perovskite scaffold, resulting in the high resistance to agglomeration and hydrocarbon coking 5 , thus maintaining the catalytic activity of nanoparticles during the long-term operation. However, the stability of P-eNs is still unsatisfactory for CO2 electrocatalysis at high voltages and has not been resolved so far. The significant current density loss has been mentioned in previous studies and is suggested to be caused by the carbon dioxide starvation at high reaction rates, coke formation by CO disproportionation/ electrochemical reduction and mass transport limitation 32,33,34 .
To address this challenge, we focus on developing a strategy to inhibit the structural evolution of the perovskite scaffold. Taking into account the promoting effect of the applied potentials on exsolution, the continuous exsolution occurs at high voltages 23 , which may cause the accumulation of B-site vacancies and the consequent phase decomposition of perovskite substrate. Therefore, we infer that the structural stability would be greatly improved after supplementing the B-site vacancies with the redox-stable cations. To achieve B-site supplement for boosting the stability of P-eNs, there are two synthesis methods to prepare the P-eNs with sufficient B-site occupation: the A-site deficiency/B-site excess and the topotactic ion exchange. However, it has been reported that pure Sr1.9Fe1.3Ni0.2Mo0.5O6-δ cannot be obtained with 5% mol Sr-site deficiency 35 , while almost all Ni elements can be exsolved from SFNM after reduction treatment at 800 o C and a 5% H2/N2 atmosphere for 2 h. In this case, less than 5% mol Sr-site deficiency cannot effectively retain the occupation of the B-site vacancies after exsolution. Therefore, we choose the ion exchange assisted exsolution method to achieve B-site supplement of perovskite scaffold of reduced SFNM.
Although the ion exchange method is not the novelty of this work and we only used the ion exchange method to enhance structure stability, thereby to facilitate stable CO2 electrolysis performance at high voltage, we have optimized the precursor deposition method for initiating this synthesis route. The simple combination of lyophilization and TIE not only achieves the B-site supplement, but also shows great application prospects in decorating the perovskite scaffold of P-eNs based on the demand of specific application.
Figs. 1c and 1d are the schematic diagrams based on the specific SFNM double perovskite structure, intended to help readers gain a better understanding on the DFT calculation and the influence of B-site supplement on the structure of the reduced SFNM perovskite scaffold. Because the similar synthetic method was used as in the Reference 26, it is reasonable that the schematic representations would be similar. We believe that combining 2D and 3D schematics in a generally recognized way is helpful for readers to understand the work of the DFT and synthetic parts.
We have revised the manuscript (L10-19, P4) accordingly, as shown below: Revised manuscript (L10-19, P4): "In this study, the promising double perovskite Sr2Fe1.3Ni0.2Mo0.5O6-δ (SFNM) was selected as a prototype example to elaborate the effects of the structure evolution of the perovskite scaffold on the stability of P-eNs 3 . Either controlling the A-site deficiency or implementing the topotactic ion exchange (TIE) is expected to be a pathway to regulate the concentration of the B-site vacancies in the reduced SFNM (Eqs. 2 and 3) 22,26,36,37 . However, the limited Sr-site deficiency (less than 5% mol) in the SFNM makes it fail to refill the B-site vacancies after the exsolution by controlling the A-site deficiency 35 . Therefore, the TIE-assisted exsolution was employed to fine-tune the B-site occupation of perovskite scaffold while promoting the formation of nanoparticles." Authors' response to Comment #2: We thank the reviewer for pointing this out. As suggested by the reviewer, all the characteristic peaks of Fig.  R11 (also Supplementary Fig. 2) have been designated in the revised manuscript. In addition, it is expected that the trace amount of secondary phase SrMoO4 will be detected in the air-sintered SFM 38 , so the corresponding PDF card has been added. From an electroneutrality viewpoint, the high oxidation state of Fe (+2/+3) and Mo (+5/+6) at the B-site of SFM would result in an unbalanced positive charge, thus leading to a limited solubility of Mo at B-site of SFM. Consequently, Mo exists as a secondary phase SrMoO4 by combing with Sr while Fe is completely dissolved in the perovskite lattice. Instead, the pure SFM can be synthesized in a reducing environment by reducing the B-site average valence 38 .

Authors' response to Comment #3:
It is true that similar host material has already been reported by other researchers. The main achievement of our work is to propose and verify the effectiveness of B-site supplement strategy for stability enhancement of the P-eNs. For this purpose, a typical P-eNs should be selected as a prototype for elucidation. Sr2Fe1.5Mo0.5O6−δ (SFM) has great potentials in solid oxide cells due to its satisfactory redox stability and conductivity under both reducing and oxidizing conditions 39,40,41 . SFM with Fe-Ni alloy nanoparticles have been regarded as the promising candidate for CO2 reduction in SOEC 3 . Like other P-eNs, the reduced Sr2Fe1.3Ni0.2Mo0.5O6−δ (SFNM) with exsolved Fe-Ni nanoparticles suffers significant degradation for CO2 electrocatalysis at high voltages, which greatly limits its further application. Therefore, SFNM was selected as a prototype example.
Comment #4. Why is there non-linearity from the voltage range of 0 to about 0.8 V in the I-V curves in Supplementary Figure 9? Also, why is the open-circuit voltage too low?
Authors' response to Comment #4: The change in the slope of the -V from the non-linearity to linearity at approximately 0.8 V separates the curve into two major cell processes. The nonlinear part (below 0.8 V) is the activation region, i.e., the SOEC mainly undergoes activation polarization and the electrolysis process is controlled by the reduction reaction on the cathode and the oxidation reaction on the anode; while the linear segment is ohmic loss region, the electrolysis process is controlled by the resistance to the flow of current through the cell when the applied potential enters into ohmic region 42,43,44,45,46,47 .
The measured open-circuit voltage (OCV) is expected to be slightly lower than the theoretical reversible voltage ( ) required for CO2 splitting. at fixed pressure and temperature can be expressed by the Nernst equation: In our experiments, the SOECs were running at 800 and 850 o C, the anode was exposed to ambient air while the cathode side was fed with pure CO2. Although the in pure CO2 environment cannot be obtained by the above Equation, value shows a decreasing trend with decreasing CO concentration. To verify the rationality of the OCVs obtained in our experiments, we tested the changes in OCV of SFNM+0.0Fe-red-GDC at 800 o C and atmosphere pressure when the cathodic atmosphere was switched from 70% CO2/30% CO to pure CO2. The measured OCV in the 70% CO2/30% CO atmosphere is 0.88 V (Fig. R12), very close to the = 0.895 calculated from the Nernst Equation ( = 0.97 at atmosphere pressure and 800 o C 48 ). The remarkably decreases to 0.08 V after switching the atmosphere from 70% CO2/30% CO to pure CO2. Based on above discussion, the measured OCVs for CO2 electrolysis in our work are all within a reasonable range calculated by the Nernst equation and similar to the OCVs reported in the literatures measured under the same conditions 49,50 .

Authors' response to Comment #6:
We thank the reviewer for the valuable comment. It is common for TGA plot to have multiple slopes in this system, and each individual slope can be ascribed to different process, such as the loss of the lattice oxygen caused by the reduction of B-site cations and exsolution of Fe-Ni alloy, as shown in the following equations 51 : Loss of lattice oxygen caused by reducing oxidation states of Fe and Ni: Loss of lattice oxygen caused by exsolution of Fe and Ni: To reveal the detailed processes responsible for the slope changes from 400 to 800 o C, the ex-situ XRD and SEM results of SFNM+0.0Fe-red and SFNM+1.2Fe-red treated at 400, 500, 600, 700, 800 o C in 5% H2/95% N2 atmosphere are provided in Figs. R13 and R14. For SFNM+0.0Fe-red, there is only a sharp weight loss between 400 and 500 o C. From the XRD results, we can see that the main peak at about 32 o shifts to the right slightly. It may be ascribed to the reduction of reducible Fe and Ni cations, which leads to the lattice expansion. The SEM results of SFNM+0.0Fe-red-500 o C-0h show that trace amount of exsolved nanoparticles has formed on the surface. It indicates that the onset temperature of nanoparticle exsolution on SFNM is below 500 o C. For SFNM+ Fe ( =0.5, 0.8, 1.2), the ramps of weight loss curves appear at slightly lower temperatures (below 400 o C), indicating that the reduction and exsolution may happen at the lower temperatures, which is consistent with the promoting effect of ion exchange on exsolution. It is also verified by the XRD results that the main peak of SFNM+1.2Fe-red shifts slightly to lower angle from 400 to 500 o C. As temperature ramping, there is a sharp weight loss between 600 and 700 o C for SFNM+ Fe ( =0.5, 0.8, 1.2). It may be ascribed to the promotion effect of TIE on nanoparticle exsolution, which results in the formation of Fe-Ni nanoparticles and loss of oxygen vacancies. The exsolved nanoparticles can be easily observed on the surfaces of SFNM+0.0Fe-red and SFNM+1.2Fe-red. Figure R13. Ex-situ XRD results of (a1-a2) SFNM+0.0Fe-red and (b1-b2) SFNM+1.2Fe-red treated at 400, 500, 600, 700, 800 o C in 5% H2/95% N2 atmosphere. Figure R14. Ex-situ SEM results of (a) SFNM+0.0Fe-red and (b) SFNM+1.2Fe-red treated at 400, 500, 600, 700, 800 o C in 5% H2/95% N2 atmosphere.
We have revised the Supplementary information (SI) (L12-14, P18 and L1-7, P19) accordingly, as shown below: Revised SI (L12-14, P18 and L1-7, P19): "The gradual weight loss below 400 o C can be ascribed to the detachment of surface-absorbed H2O molecule and decomposition of nitrate 14 . Further raising the temperature induces a continuous decline in sample weight for all samples, which was attributed to the formation of oxygen vacancy caused by the reduction of B-site cations and exsolution of Fe-Ni alloy, as shown in the following equations 51 : Loss of lattice oxygen caused by reduced oxidation states of Fe and Ni: Loss of lattice oxygen caused by exsolution of Fe and Ni: Are there any additional supporting data (e.g., X-ray absorption fine structure (XAFS) measurement) to further support the XPS data in Supplementary figure 13?
Authors' response to Comment #7: Thanks to the reviewer for the suggestion. As suggested by the reviewer, the Fe and Ni K-edge X-ray absorption near edge structure (XANES) spectra among air-sintered SFNM-oxi, SFNM+0.0Fe-red and SFNM+1.2Fe-red have been provided in Fig. R15 (also Supplementary Fig. 17) 52,53 . The shifts of pre-edge of Fe and Ni K-edge for SFNM+0.0Fe-red and SFNM+1.2Fe-red to lower energy with respect to that for SFNM-oxi are associated with the decrease of valence states of Fe and Ni after reduction 54,55 . It is clear that the valence states of Ni are almost identical in SFNM+0.0Fe-red and SFNM+1.2Fe-red, while the Fe valence state for SFNM+1.2Fe-red is apparently lower, which is consistent with the Fe and Ni 2p XPS spectra. It suggests that the amount of the Ni exsolution has almost reached the peak value via conventional exsolution due to its lower co-segregation energy. Figure 17). Fe and Ni K-edge XANES spectra of sintered SFNM, SFNM+0.0Fered and SFNM+1.2Fe-red.

Figure R15 (also Supplementary
We have revised the SI (L5-14, P23) accordingly, as shown below: Revised SI (L5-14, P23): "The Fe and Ni K-edge X-ray absorption near edge structure (XANES) spectra among airsintered SFNM-oxi, SFNM+0.0Fe-red and SFNM+1.2Fe-red are provided in Supplementary Fig. 17 52,53 . The shifts of pre-edge of Fe and Ni K-edge for SFNM+0.0Fe-red and SFNM+1.2Fe-red to lower energy with respect to that for SFNM-oxi are associated with the decrease of valence states of Fe and Ni after the reduction 54 , 55 . It is clear that the valence states of Ni are almost identical in SFNM+0.0Fe-red and SFNM+1.2Fe-red, while the Fe valence state for SFNM+1.2Fe-red is apparently lower, which is consistent with the Fe and Ni 2p XPS spectra. It suggests that the amount of the Ni exsolution has almost reached the peak value via conventional exsolution due to its lower co-segregation energy." Comment #8. In my opinion, this work did not follow the Nature Communications format. Is the font for this work "times" or "times new roman"?
Authors' response to Comment #8: The format of the revised manuscript follows the requirement of Nature Communications and the font is "Calibri".
Comment #9. Why did the authors use electrode-GDC composite for SOEC measurements? Since the GDC could significantly affect the exsolution and/or TIE properties, the authors should compare the electrochemical performance measurements with only electrode material (SOEC cathode).

Authors' response to Comment #9:
We thank the reviewer for pointing this out. For electrodes in SOEC, the perovskite-based materials are usually mixed with a pure oxygen ion conductor to prolong the triple-phase boundaries and reaction sites, thereby resulting in a higher electrocatalytic performance than their pure perovskite cathode counterpart 56,57 . The common pure ion conductors include gadolinium-doped ceria (GDC) and samarium-doped ceria (SDC). The CO2 electrocatalysis process can be described using Kröger-Vink notation: The reduction of GDC ( 4+ → 3+ ) in the composite electrodes under cathodic polarization would provide additional electrical conduction paths 58,59 .
Most importantly, since there is a GDC barrier layer between the electrode material and YSZ electrolyte in the cell assembly (LSCF-GDC|GDC|YSZ|GDC|SFNM+ Fe-red-GDC), the composite cathodes are more chemically compatible with the GDC buffer layer. It would reduce the difference of thermal expansion coefficient and interfacial resistance between cathode and GDC buffer layer 57 . For the same reason, LSCF-GDC composite was selected as anode electrode. Furthermore, for the material preparation and cell assembly, the SFNM+ Fe-red ( =0.0, 0.5, 0.8, 1.2) have been prepared by pre-reduction in 5% H2/95% N2 before mixing with GDC, so it would not affect the conventional and TIE-assisted exsolution processes.
Comment #10. In page 7 line 15, the authors stated as "X-ray diffraction (XRD) analysis confirms that secondary Fe-Ni alloy phases have emerged …". In general, secondary phase could also imply unnecessary or impurity phase. However, the Fe-Ni alloy exsolution would positively affect the electro-catalytic properties, as also listed in this work.
Authors' response to Comment #10: Thanks to the reviewer for this valuable suggestion, the "secondary Fe-Ni alloy phases" has been corrected to "Fe-Ni alloy phases" on L15, P8 of the revised manuscript.
Comment #11. What is the possible reason for ultrasound (ultrasonication) and freeze-drying Fe nitrate nonahydrate in the host material (Supplementary figure 1)? In addition, how much weight percent of Fe guest did the authors immerse in the SFMN material to promote TIE phenomenon?
Authors' response to Comment #11: The water solubility of ferric nitrate and the high temperature decomposition of nitrate after heat treatment to 800 o C are the main reasons why we choose it as the Fe source to be loaded on the surface of perovskite scaffold. The ultrasound treatment followed by freeze-drying treatment allows the external Fe source to be uniformly deposited on the perovskite substrates. The principle of freeze-drying is to utilize the sublimation of objects. In a vacuum environment, the frozen water can directly become gas without going through the intermediate liquid phase, leaving the Fe source on the perovskite substrates.
The Fe loading amount on the surface of perovskite substrates has been defined on L14, P8 in the manuscript "a series of SFNM+ Fe-red samples ( =0.0, 0.5, 0.8, 1.2, which refer to the molar ratios of guest Fe to host Ni)". After determining the weight of SFNM powder, the weight ratio of ferric nitrate for preparing SFNM+ Fe-red samples can be calculated by simple unit conversion (mole to gram). For instance, if using 0.3 g SFNM powder to prepare SFNM+ Fe-red ( =0.0, 0.5, 0.8, 1.2), the weight of ferric nitrate required is shown in Table R5:  Figure 2, there seems to be secondary phases (NiO, NiFe2O4, and SrMoO4) after re-oxidation in air. The formation of secondary phases could negatively affect the electrochemical properties. Moreover, the reason for conducting re-oxidation in air is not clearly explained in this work.

Authors' response to Comment #12:
Thanks to the reviewer for raising the question. The reoxidation experiments is to further verify the reoccupation of B-site by the guest Fe for SFNM+ Fe-red samples ( =0.5, 0.8, 1.2).
It has been reported that the exsolution of nanoparticles from perovskites is reversible/partially reversible 24,60 , i.e., the exsolved nanoparticles could completely/partially dissolve into the lattice of parent perovskite by the re-oxidation treatment. To verify this, the XRD analyses were conducted. The results show that the additional NiO peaks have clearly emerged in re-oxidized SFNM+0.0Fe-red, indicating that the redissolution of the exsolved Fe-Ni nanoparticles on the SFNM is partially reversible. Moreover, Fe atoms in Fe-Ni nanoparticles preferentially dissolve into perovskite compared to Ni atoms, which is in accord with the fact that Ni exsolves more favourably than Fe. However, NiFe2O4 or SrMoO4 can be detected on re-oxidized SFNM+ Fe-red ( =0.5, 0.8, 1.2). The appearance of SrMoO4 indicates that the B-sites have been almost occupied by the high-valence Fe and Mo 38 , the exsolved Fe and Ni on surface are prone to self-assembly into binary oxide NiFe2O4 in air.
Comment #13. Why did the authors conduct reduction at 5% H2/N2 balance instead of real experimental conditions for SOEC cathode (Figure 2)?

Authors' response to Comment #13:
The pre-reduction treatment of sintered perovskites aims to produce the highly active metal nanoparticles on the surface by exposing the perovskite to a reducing atmosphere at elevated temperatures, thus enabling it to have a higher initial reactivity for CO2 electrocatalysis. Fig. R16 shows the surface morphologies of exsolved nanoparticles on SFNM+0.0Fe-red and SFNM+1.2Fe-red after heating at 800 o C for 2h in 70% CO2/30% CO atmosphere. The populations of the exsolved nanoparticles on both samples treated in 70% CO2/30% CO environment are far less than those treated in 5% H2/95% N2. It suggests that the 5% H2/95% N2 atmosphere is more potent for nanoparticle exsolution than 70% CO2/30% CO environment 34 , which delivers a higher catalytic activity for CO2 reduction. This stark difference in the surface features of reduced perovskites may be dominated by the different oxygen partial pressure and surface defect reactions in different reducing atmosphere 17,34 . The exsolution process of Fe and Ni from SFNM perovskite in the two types of atmospheres (H2/N2 and CO/CO2) can be expressed by the following equations: Exsolution at H2/N2 atmosphere: Exsolution at CO/CO2 atmosphere: The exsolved Fe-Ni alloy nanoparticles have been regarded as the efficient CO2 adsorption sites on the surface of perovskites. Consequently, it would block the binding of CO molecules with the surface of perovskites, which may inhibit the continuous exsolution 34 , as shown from the above reaction process. Therefore, the H2/N2 reducing atmospheres have been widely applied to produce the P-eNs in the literatures and our work.
Comment #14. There seems to be SrO segregation in Figure 3d. In general, the A-site segregation hampers the exsolution capability and could also negatively affect the electro-catalytic properties.

Authors' response to Comment #14:
We agree that the A-site segregation would hamper the exsolution capability and could also negatively affect the electro-catalytic properties. As seen from Figure 3d, SrO segregation occurs at around the nanoparticles of the two samples, and the undesirable Sr-segregation of SFNM+1.2Fe-red is less severe than that of SFNM+0.0Fered. It is mainly due to the B-site supplement, as explained by Eqs. 6 and 7 on P14-15 in the revised manuscript. Figure 4c. In the case of SFNM+1.2Fe-red-GDC (Figure 4c), there seems to be about 30% degradation after 100-hour stability test.

Authors' response to Comment #15:
We agree with the reviewer. The observable degradation is mainly ascribed to the harsh testing conditions (1.6 V and 850 o C), which cause the degradation of not only cathode, but also the anode and electrolyte. As mentioned on L13-18, P19 in the revised manuscript, the high cathode potential and temperature are applied to intentionally accelerate the degradation of SOECs to help us understand the cathode evolution and further shed lights on the degradation mechanisms. The SEM results of cathode surface morphologies on the SFNM+0.0Fe-red-GDC and SFNM+1.2Fe-red-GDC as well as the cross-section of electrolysis cells after the longterm stability are characterized, as shown in Figs. R17 and R18 (also Figs. 5d-e and Supplementary Fig. 24). The obvious morphology changes and coarsening of exsolved particles can be observed on both cathodes, and the coarsening of surface nanoparticles on SFNM+1.2Fe-red-GDC is less severe than that on SFNM+0.0Fe-red-GDC (Fig. R17). Furthermore, the significant grain coarsening of YSZ electrolyte and the delamination of anode from the electrolyte can be observed on both cells after the long-term stability tests and both factors also contributed to the degradation (Fig. R18). Figures 5d-e). SEM images of cathode surface microstructures of (a) SFNM+0.0Fe-red-GDC and (b) SFNM+1.2Fe-red-GDC after long-term stability at 1.6 V and 850 o C. Figure 24). SEM images of cross section of (a) SFNM+0.0Fe-red-GDC and (b) SFNM+1.2Fe-red-GDC after long-term stability at 1.6 V and 850 o C.

Figure R18 (also Supplementary
We have revised the manuscript (L1-5, P21, L14-18, P25) and SI (L5-14, P31) accordingly, as shown below: Revised manuscript (L1-5, P21): "After the long-term stability tests, the significant morphology changes and coarsening of exsolved particles can be observed on the both cathode surfaces, and the coarsening of surface nanoparticles on SFNM+1.2Fe-red-GDC is less severe than that on SFNM+0.0Fe-red-GDC (Figs. 5d-e)." (L14-18, P25): "In addition, the deteriorations of the electrolyte and the anode under the harsh conditions seem to be inevitable ( Supplementary Fig. 24), which are also responsible for the significant attenuation of the current density during CO2 electrocatalysis at 1.6 V and 850 o C." Revised SI (L5-14, P31): "The degradation of the electrolyte and the anode under the harsh conditions are analyzed 61 . Supplementary Fig. 24 shows the SEM images of the cross section of SFNM+0.0Fe-red-GDC and SFNM+1.2Fe-red-GDC after the long-term stability tests at 1.6 V and 850 o C. The significant grain coarsening of YSZ electrolyte can be observed near the anode|GDC|YSZ interfaces of both cells, and this grain growth propagates along the electrolyte and progresses towards the cathode side. Furthermore, the pore formation near the anode|GDC|YSZ interfaces and even the delamination of anode from the electrolyte can be observed, which is supposed to originate from the high oxygen partial pressures formed at the interfaces during the longterm CO2 electrolysis at high voltage." Comment #17. In the abstract part, the authors stated as "In this study, we reveal that the formation of Bsite vacancies in perovskite scaffold is the major contributor to the degradation of P-eNs; …". Instead of Bsite vacancies in perovskite scaffold, A-site segregation and/or the formation of impurity phases (SrCO3 and SrMoO4) in Figure 4e appears to be much significant factor for the degradation properties. This part should be clearly elucidated.
We thank the reviewer for pointing this out. From our work, we conclude that the structural decomposition of perovskite scaffold is caused by the accumulation of B-site vacancies during the CO2 electrolysis at high voltages. During exsolution, the reducible B-site ions migrate to the surface to form nanoparticles, accompanied by the formation of B-site vacancies. When a small number of B-site reducible cations exsolve out of the perovskite lattice, the concentration of B-site vacancies is not enough to reach the threshold of destructive structural decomposition since perovskite has a certain structural tolerance. When the created B-site vacancies exceeds the threshold, the perovskite partially decomposes into Sr-based derivatives on the surface. Therefore, the Asite segregation and the formation of impurity phases (SrCO3 and SrMoO4) during exsolution are the significant factor for the degradation properties. As discussed in the revised manuscript (L3-4, P15): "The Sr-rich surface would reorganize into SrCO3 at a high CO2 concentration, which has a detrimental effect on the surface catalytic activity for CO2 reduction 62 ." and the revised manuscript (L16-18, P21): "The emergence of Sr-containing impurities on both spectra confirms the structure decomposition and surface reconstruction on SFNM+0.0Fered and SFNM+1.2Fe-red." and the revised manuscript (L5-8, P22): "As shown in Fig. 6b, only the Raman feature peaks of SrCO3 and SrMoO4 were detected but no peaks of carbon, which confirms that the perovskite scaffold of SFNM+0.0Fe-red primarily underwent the bulk structural decomposition before the exposure of newly grown nanoparticles and carbon deposition." Comment #18. The current density for SFNM+0.0Fe-red-GDC at 1.8 V much decreases just after 15 minutes. Hence, the production rate of CO and the Faradaic efficiency of CO should be much different between 60 min and 75 min.

Authors' response to Comment #18:
Thanks for this question. The Faradaic efficiency of CO is calculated by the following equation 63,32 : where ( ) = Gas flow rate measured by a flow meter at the exit of the cell at room temperature and under ambient pressure.
( %) = Volume concentration of CO in the exhaust gas from the cell (obtained by gas chromatography).

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As shown in the equation, depends on the ratio of ( %) to ( ) . The current density for SFNM+0.0Fe-red-GDC at 1.8 V drops remarkably after 15 minutes, but in the meantime, the ( %) decreases at the same degree accordingly. Therefore, the value of would not change significantly.
The following equation is used to calculate the production rate of CO 64  = the geometric area of the electrode (cm 2 ).
The production rate of CO is proportional to the which is the integral of cell current versus time and is an average value over 15 min. Fig. 5b in the revised manuscript shows the comparison of the average CO productivity of short-term stability tests at different potentials. It can be seen that the current density for SFNM+0.0Fe-red-GDC at 1.8 V drops remarkably during the stability test, the total charge passing through SFNM+0.0Fe-red-GDC at 1.8 V is much smaller than that at 1.6 V during the 15 min of test. Therefore, the CO productivity at 1.8 V is smaller than that at 1.6 V during the 15 min of test, as shown in Fig. 5b. In addition, the instantaneous production rate of CO should decrease with decreasing the current density. Hence, the CO productivity for SFNM+0.0Fe-red-GDC at 1.8 V decreases significantly from 60 min to 75 min. Comment #19. How did the authors calculate the Faradaic efficiency of CO and the production rate of CO via gas chromatography measurements? I couldn't find clear explanations and equations on how the authors calculated the Faradaic efficiency of CO and the production rate of CO.

Authors' response to Comment #19:
We thank the reviewer for this comment. We have added the equations of the Faradaic efficiency of CO and the production rate of CO to the "Method section" in the revised manuscript (P33-34). Comment #20. Did the authors calculate the Gibbs free energy for Fe-Ni alloy formation? Furthermore, is there any Gibbs free energy calculations on whether the Fe-Ni alloy formation is favorable at the surface or the bulk? This part should be also explained in this paper since the title in this work is "B-site supplement mechanism: A new approach to boosting stability of perovskites with exsolved nanoparticles".

Authors' response to Comment #20:
Thanks for the suggestion. The Gibbs free energy for Fe-Ni alloy formation is not calculated in the manuscript because we can know that the Fe-Ni alloy can be formed from the XRD and TEM-EDS results in this work. The formation location of alloy in perovskites is a very interesting topic. It is a challenging work to unveil the formation location of alloy nanoparticles by DFT calculation since it is not easy to define the accurate configuration of alloy formed in the perovskite bulk from the perspective of energy. Therefore, it may not be possible to correctly judge the preferred sites of alloy formation by DFT method. To the best of our knowledge, only the work of Kwon et al. was to explore the bulk Co-Ni alloy formation using DFT calculation 65 . They defined the configuration in which Co and Ni share the same oxygen vacancy in bulk, but this definition does not really represent the actual situation.
Alternatively, it has been recently reported that the surface exsolved nanoparticles from (La0.3Ca0.7)0.95Fe0.7Cr0.25Ni0.05O3-δ grown in a H2-N2 atmosphere exhibit a compositional progression from an initially Ni-rich phase to an Fe-rich phase during the heat treatment 34 . It indicates that the formation of surface Fe-Ni alloys have undergone the process of diffusion of Fe and Ni from the bulk to the surface followed by alloying at surface. In addition, two other typical studies on the metallic nanoparticle nucleation mechanism demonstrated that the exsolved metallic nanoparticles would nucleate at surface or subsurface by in-situ TEM, atomic force microscopy (AFM) and strain field modeling 25,31 . These two metallic nanoparticle formation scenarios are more acceptable compared to the bulk nanoparticle formation due to the high strain from the surrounding lattice and high diffusion energy barrier of the nanoparticles formed in the perovskite bulk. Besides, Kousi et al. reported the endogenous nanoparticles nucleation in perovskite bulk by selecting a titanate perovskite with relatively low cation transport 28 . They proposed that the design principle of bulk nucleation is to make the diffusivity of B-site reducible cations in the bulk as low as possible, so that the B-site reducible cations favour local particle nucleation rather than transport to the surface followed nucleation. All these evidences demonstrate that the favourable location for the formation of exsolved metallic nanoparticle is at surface rather than in bulk. Furthermore, the main focus of our work is to enhance the stability of P-eNs rather than to investigate the formation location of alloy exsolution Thus, our goal of this work is only focused on calculating the co-segregation energy of different B-site cations of SFNM to judge the external B-site supplemental element. Comment #21. The overall thesis design and/or English proficiency should be much complemented in this work.

Authors' response to Comment #21:
We appreciate the reviewer's comment and have carefully revised the manuscript to improve its readability.

Reviewer #3 (Remarks to the Author):
This work proposed B site supplement strategy to compensate the occurrence of B site vacancies caused by exsolution process before and during the operation, especially at high reduction potential. Mostly, nanoparticle exsolution of B site cations in perovskite lattice leaves many B site vacancies and causes detrimental A site segregation, leading to the losses of B (n+1) -O-B n+ pathway. In this work, authors have insisted that the guest Fe metal coated on Sr2Fe1.3Ni0.2Mo0.5O6-δ (SFNM) would fill the B site vacancies by topotactic ion exchange (TIE) mechanism and mitigate the Sr segregation on the surface, thus enhancing structure stability and facilitating stable CO2 electrolysis performance at high voltage. Several characterizations such as XRD, SEM, TEM, and elemental mappings were performed to verify the role of guest Fe, filling the B site vacancies during exsolution process and suppressing Sr segregation on perovskite surface. In potential step chronoamperometry, SFNM without guest Fe (SFNM+0.0Fe-red) showed significant decay of current density (j), whereas SFNM with guest Fe metals (SFNM+1.2Fe-red) showed stable j profiles, especially under high operating voltage (1.8V). In this work, the results are noteworthy and would be useful to develop a highly active CO2 electrolysis cathode catalyst with superior stability in SOEC system. Authors have provided detailed mechanisms of perovskite structure evolution during the exsolution and TIE-assisted exsolution, and most of conclusions are supported by the results. It might be suitable to be published in this journal after addressing below issues and comments:

Authors' response:
We thank Reviewer #3 for the highly positive comments and recommendation for publication. Our detailed responses to each comment are as follows.
Comment #1. In page 3~4, authors have mentioned that the exsolution process on the perovskite leaves many B site vacancies within perovskite bulk and thus A site segregation on the surface as below equation: ABO_(3-δ) → A_(1-α')B_(1-α)O_(3-δ) + αB + α'AO However, it is well known that these defects (i.e., B_(1-α) and AO) can be relieved by controlling A site nonstoichiometry on perovskite lattice as followed: A-site deficient design has been adopted in numerous other researches of SOEC CO2 cathodes, especially in [D. Neagu et al., Nat. Chem., 2013, 5, 11, 916-923]. It might be helpful to mention on your manuscript that what is the strong point of your strategy to coat additional guest Fe metals on perovskite surface to induce TIE-assisted exsolution compared to the one, which controls the A site non-stoichiometry of perovskite oxide.
Thanks to the reviewer for this valuable suggestion. Indeed, as mentioned by the reviewer, A-site deficiency/Bsite excess is an alternative way to reduce the concentration of the B-site vacancies of perovskite scaffold in addition to the ion exchange. These two schemes had been considered when we proposed the B-site supplement mechanism. However, the adjustable range of A-site deficiency in Sr2Fe1.3Ni0.2Mo0.5O6-δ (SFNM) is relatively limited. It has been reported that pure Sr1.9Fe1.3Ni0.2Mo0.5O6-δ cannot be obtained with 5% mol Sr-site deficiency 35 . Moreover, almost all Ni elements can be exsolved from SFNM after reduction treatment at 800 o C and a 5% H2/N2 atmosphere for 2 h. In this case, less than 5% mol Sr-site deficiency cannot fully retain the occupation of the B-site vacancies after exsolution. Therefore, we choose the ion exchange assisted exsolution method to achieve B-site supplement of perovskite scaffold of reduced SFNM. As suggested by the reviewer, we have added the following paragraph on L10-19, P4 in the revised manuscript to explain the strong point of the strategy to coat additional guest Fe sources on the perovskite surface to induce TIE-assisted exsolution compared to the one controlling the A site non-stoichiometry of perovskite oxide.
Revised manuscript (L10-19, P4) : "In this study, the promising double perovskite Sr2Fe1.3Ni0.2Mo0.5O6-δ (SFNM) was selected as a prototype example to elaborate the effects of the structure evolution of the perovskite scaffold on the stability of P-eNs 3 . Either controlling the A-site deficiency or implementing the topotactic ion exchange (TIE) is expected to be a pathway to regulate the concentration of the B-site vacancies in the reduced SFNM (Eqs. 2 and 3) 22,26,36,37 . However, the limited Sr-site deficiency (less than 5% mol) in the SFNM makes it fail to refill the B-site vacancies after the exsolution by controlling the A-site deficiency 35 . Therefore, the TIE-assisted exsolution was employed to fine-tune the B-site occupation of perovskite scaffold while promoting the formation of nanoparticles." Comment #2. In page 6, co-segregation energies of Ni, Fe and Mo at B site of SFNM were calculated by DFT method. From the calculation results, it seems to be reasonable to choose Ni as a guest ions to induce vigorous TIE-assisted exsolution because of its relatively low segregation energies (Ni: -1.46 eV, Fe: -1.06 eV, Mo: 2.26 eV). However, authors have chosen Fe as guest metals due to its relatively high redox stability (in page 5). Is there any additional calculation results that can support the choice of Fe as a guest ions?
Authors' response to Comment #2: The following factors are mainly considered when choosing the B-site supplement agent: 1) the doping availability of deposited cations into SFNM perovskite lattice; 2) the exchangeability of deposited ions with the easily reducible Fe and Ni in the parent SFNM lattice.
Firstly, the Sr2Fe1.3Ni0.2Mo0.5O6-δ (SFNM) can be seen as the Sr2Fe1.5Mo0.5O6-δ (SFM) with the partial Ni substitution at Fe-site, which causes the structural instability of perovskite scaffold and yields the highly active Fe-Ni alloy after the reduction. Therefore, the redox-stable Fe, instead of the highly reducing Ni, is expected to supplement into B-site vacancies, thus leading to a relatively robust perovskite scaffold.
We clearly know that Fe and Ni can be exsolved from SFNM, leaving the Fe-and Ni-site vacancies, so the external Fe source must be able to exchange with the host Fe and Ni to achieve B-site full occupation. The exchangeability is determined by comparing the co-segregation energies of B-site elements 26,37 . The elements with lower cosegregation energy tend to exsolve on the surface, in turn, the elements with higher co-segregation energy tend to dissolve into the perovskite lattice. The calculated co-segregation energy trend of B-site elements is Ni>Fe>Mo. Thus, the deposited Fe can exchange with host Ni and Fe but not the host Mo, resulting in the supplement of Fe and Ni-site vacancies of reduced SFNM with foreign Fe. However, the deposited Ni can only exchange with the host Ni, leaving a large number of Fe-site vacancies in SFNM. Therefore, the relatively redoxstable Fe is selected to supplement into the B-site vacancies during exsolution.

Comment #3.
Authors have insisted that TIE-assisted exsolution would decrease B site vacancies that could be produced during exsolution, thus help to maintain the electrical conductivity of bulk perovskite. XPS analysis, which showed more Fe 2+ -Fe 3+ and Mo 5+ -Mo 6+ charge pairs in SFNM+1.2Fe-red, was conducted to support the opinions. Is it possible to measure electrical conductivities of SFNM+0.0Fe-red and SFNM+1.2Fered under reducing and operating atmosphere?

Authors' response to Comment #3:
We thank the reviewer for this valuable suggestion. As suggested by the reviewer, the temperature-dependent electrical conductivities of SFNM+0.0Fe-red and SFNM+1.2Fe-red in the 50% CO2/50% CO and 5% H2/95% Ar atmospheres were measured and the results are shown in Fig. R19 (also Fig. 4d and Supplementary Fig. 18). In the 50% CO2/50% CO atmosphere, the electrical conductivities of both SFNM+0.0Fe-red and SFNM+1.2Fe-red increase as the temperature increases, reaching 16.5 and 24.6 S cm -1 at 850 o C, respectively (Fig. 4d). The results clearly show that SFNM+1.2Fe-red has the higher conductivity in the prospective operation atmosphere of CO2 electrolysis. The electrical conductivities of SFNM+0.0Fe-red and SFNM+1.2Fe-red in the 5% H2/95% Ar atmosphere at 850 o C reach 26.0 and 48.2 S cm -1 , respectively, which are higher than those in the 50% CO2/50% CO atmosphere. It can be ascribed to the contribution of the more exsolved Fe-Ni alloy nanoparticles formed under the more reducing condition in addition to the conduction pathways in perovskite scaffold 66 . A large number of metallic nanoparticles on SFNM+1.2Fe-red leads to the slightly reduced conductivity as the temperature increases in the 5% H2/95% Ar atmosphere. Figure 4d and Supplementary Figure 18). Temperature-dependent electrical conductivities of SFNM+0.0Fe-red and SFNM+1.2Fe-red under 50% CO2/50% CO and 5% H2/95% Ar atmospheres.

Figure R19 (also
We have revised the manuscript (L16-19, P17 and L1-4, P18) and SI (L5-12, P24) accordingly, as shown below: Revised manuscript (L16-19, P17 and L1-4, P18): "The temperature-dependent electrical conductivity results show that SFNM+1.2Fe-red performs better than SFNM+0.0Fe-red in the 50% CO2/50% CO and 5% H2/95% Ar atmospheres ( Fig. 4d and Supplementary Fig. 18). In the 50% CO2/50% CO atmosphere, the electrical conductivities of SFNM+0.0Fe-red and SFNM+1.2Fe-red increase as the temperature increases, reaching 16.5 and 24.6 S cm -1 at 850 o C, respectively (Fig. 4d). SFNM+1.2Fe-red shows the higher conductivity in the prospective operation atmosphere of CO2 electrolysis." Revised SI (L5-12, P24): "As depicted in Supplementary Fig. 18, the electrical conductivities of SFNM+0.0Fe-red and SFNM+1.2Fe-red in the 5% H2/95% Ar atmosphere at 850 o C reach 26.0 and 48.2 S cm -1 , respectively, which are higher than those in the 50% CO2/50% CO atmosphere. It can be ascribed to the contribution of the more exsolved Fe-Ni alloy nanoparticles formed under the more reducing condition in addition to the conduction pathways in perovskite scaffold 66 . A large number of metallic nanoparticles on SFNM+1.2Fe-red leads to the slightly reduced conductivity as the temperature increases in the 5% H2/95% Ar atmosphere." Comment #4. In Figure 3, two electrodes, SFNM+0.0Fe-red and SFNM+1.2Fe-red, have shown different behaviors in ohmic resistance when the external voltage is applied. It would be better to explain the change of ohmic resistance under applied voltage conditions for each electrodes.

38
Authors' response to Comment #4: We thank the reviewer for this suggestion. As suggested, the changes in the ohmic resistance (Rs) with the applied voltages have been analyzed in the revised SI (P12-13), as shown below: Revised SI (P12-13): "From 1.0 to 1.2 V, the cathode reaction rate accelerates as the voltage increases, leading to an increase of the conductivity, i.e., the oxygen ion transfer resistance is reduced at cathode side and so is RS. However, the RS values of both cells increase as the applied voltage increases to 1.4 V, which may be ascribed to the increased resistance at the anode/electrolyte interface. Because the oxygen ions produced at the anode/electrolyte interface gradually accumulate when the cathode reaction kinetics speeds up, this accumulation would hinder the transfer of oxygen ions. This explains why the RS of the SFNM+1.2Fe-red-GDC with superior oxygen ion conduction at cathode side increases by a larger magnitude than that of SFNM+0.0Fered-GDC. As the applied voltage further increases to 1.6-1.8 V (above the thermally neutral voltage EH=1.46 V at 850 o C 67,68 ), the net heat is produced because the entropic heat consumption rate becomes slower than the production rate of the irreversible heat (due to activation, ohmic, and mass transport losses in the electrolyzer 67 ). This net heat would cause the promoted transport of oxygen ions, consequently leading to a greatly reduced RS." Thanks to the reviewer for this constructive suggestion and important information. Performance comparison tables are helpful for highlighting the novelty of our work. We have cited all the references suggested by the reviewer in our revised manuscript. As discussed in our response to Reviewer #1, Comment #1A, the Nyquist plots and the fitted RS, Rp (equivalent circuit model LR(QHRH)(QLRL)) of SFNM+1.2Fe-red-GDC at 800 o C are provided. The current density, polarization resistance, CO productivity and short-term stability of SFNM+1.2Fered-GDC in our work are compared with other state-of-art P-eNs based SOECs. Please refer our response to Comment 1A, Reviewer #1 for more details.
Comment #6. What is pre-reduction conditions of SFNM+1.2-red electrode before CO2 electrolysis test?

Authors' response to Comment #6:
Sorry for having missed this information. All the pre-reduction treatments for SFNM+ Fe-red-GDC ( =0.0, 0.5, 0.8) cathode before CO2 electrocatalysis are the same. We have revised the manuscript (L11-15, P32), as shown below: Revised manuscript (L11-15, P32): "The 5% H2/N2 reducing gas flow is continuously pumped into the SFNM+ Fered-GDC ( =0.0, 0.5, 0.8, 1,2) cathode compartment during the temperature ramping of SOEC. The temperature is maintained at 800 o C for 2 h to complete the further reduction and exsolution of SFNM+ Fe-red-GDC."  As we have summarized in the first round of point-by-point response to the comment #1, the novelty of this work is not just focused on the ion exchange method. Compared with the innovation of the specific synthetic methods that the reviewer is concerned about, our work focuses more on the creative ideas and solutions to address the stability issues of P-eNs. As far as we know, as of this writing, the concept of B-site supplement strategy has not been reported in any of the previous studies.
To make the answer simple and straightforward, we would like to re-summarize the significance of our work: Firstly, the robust structure of the perovskite scaffold is identified as the main factor to enhance the stability of P-eNs, which has been neglected in the previous literatures when investigating the stability of P-eNs 1,2,3 . In this work, on the other hand, the B-site supplement strategy is proposed to enhance the structural stability of perovskite substrate, consequently leading to the higher stability of SFNM-based P-eNs. The degradation mechanism of conventional P-eNs and the role of B-site supplement strategy in boosting the stability of P-eNs have both been revealed by the intentionally accelerated degradation tests (ADT) of SFNM-based P-eNs with and without B-site supplement at high voltages. The outputs from this work can help to broaden the strategic thinking from the perspective of exsolved nanoparticles and the surface side reactions to the perovskite support in order to improve the stability of P-eNs. This bears great significance for developing the advanced P-eNs with high catalytic activity and stability for CO2 reduction and more broadly, for other energy storage and conversion systems.
Comment #2. The authors explained that the Fe source is not fully incorporated as B-site of perovskite oxides for SFM. This also indicates that B-site supplementary mechanism or TIE process can't be fully performed because of SrMoO4 formation.

Authors' response to Comment #2:
We would like to clarify the B-site supplement and re-oxidation treatment on SFNM+ Fe-red in this work.
Firstly, the surface-deposited Fe source is not only involved in the B-site supplement of perovskite scaffold, but also involved in the formation of the surface nanoparticles.
Secondly, the secondary SrMoO4 phase is not formed during the ion exchange assisted exsolution, but is formed after the further reoxidation treatment, because it has been reported that the secondary phase SrMoO4 would form in air-sintered Sr2Fe1.5Mo0.5O6-δ (SFM) 4 , 5 . From an electroneutrality viewpoint, the high oxidation state of Fe (+2/+3) and Mo (+5/+6) at the B-site of SFM would result in an unbalanced positive charge, thus leading to a limited solubility of Mo at B-site of SFM. Consequently, Mo exists as a secondary phase SrMoO4 by combing with Sr while Fe is completely dissolved in the perovskite lattice. To verify the extent to which the B-site of SFNM+ Fered ( =0.5, 0.8, 1.2) was incorporated by the external Fe source, the SFNM+ Fe-red ( =0.0, 0.5, 0.8, 1.2) was subjected to the reoxidation treatment. It demonstrates that SFNM+ Fe-red has been sufficiently filled with foreign Fe and has formed a structure close to SFM when SrMoO4 forms after the reoxidation treatment of SFNM+ Fe-red. Fig.2b clearly shows that the SrMoO4 forms during the reoxidation process of SFNM+0.8Fe-red and SFNM+1.2Fe-red, which in turn confirms that the B-site supplementary mechanism works out well on the SFNM+0.8Fe-red and SFNM+1.2Fe-red. Comment #3. The OCV value seems to be about ~0.88 V at 70% CO2/30% CO condition and ~0.1 V at pure CO2 condition (Fig. R12). However, the OCV value in Fig. 3a and Supplementary Fig. 9 seems to be much low considering the experimental conditions for SOEC measurements (70% CO2/30% CO condition), hence much different with other reported literatures regarding SOEC measurements.

Authors' response to Comment #3:
We did not utilize 70% CO2/30% CO condition in this work. For Fig. 3a and Supplementary Fig. 9, the pure CO2 was fed to the cathodic side during the CO2 electrocatalysis process, as shown on L17, P32 in the Methods part of the revised manuscript. Therefore, the OCV values shown in Fig. 3a and Supplementary Fig. 9 should be close to 0.1 V and are in consistence with those from previous work performed with pure CO2 supply. Please refer to Fig. 5a in Reference 6 published in Nature communications, Fig. 4a in Reference 7 published in Advanced Energy Materials and Fig. 10a in Reference 8 published in Journal of Materials Chemistry A. Comment #4. The shape of the XAFS figure (Supplementary Fig. 17) is too weird in my position since the characteristic line for both Fe K-edge and Ni K-edge is not in coincidence with other previously reported papers. In addition, the references for XAFS measurements is not given in this work (e.g., Fe foil, Fe2O3, NiO, etc…). Furthermore, the fourier-transformed k3-weighted plot for the XAFS measurements is not plotted in this work.

Authors' response to Comment #4:
The specific references mentioned as "other previously reported papers" are not included in the reviewer's Comment. Thus, we reviewed the literatures reporting the Fe K-edge and Ni K-edge X-ray absorption near-edge structure (XANES) spectra for the similar materials. We have found that the shape of Fe K-edge and Ni K-edge XANES spectra in our work are in consistence with those in the previous literatures 5, [9][10][11][12][13] , please refer to the Fe K-edge data (Fig. 4) in Reference 12 and the Ni K-edge data (Fig. 3) in Reference 13, both published in Nature communications.
Per the reviewer's comments, the XANES data of the references (Fe foil, FeO, Fe2O3, Ni foil, NiO) have been added (Figs. R3a-b). The Fe K-edge and Ni K-edge extended X-ray absorption fine structure (EXAFS) data are also presented (Figs. R3c-d). From the Fe K-edge EXAFS spectra (Fig. R3c), it can be seen that the overlapping area between the peak of SFNM+1.2Fe-red and the peak of Fe foil is much larger than that between SFNM+0.0Fe-red and Fe foil, which is consistent with the XPS and XANES data. This can be ascribed to the participation of external Fe sources in the formation of Fe-Ni alloy nanoparticles. As seen from the Ni K-edge EXAFS spectra (Fig. R3d), the valence state of Ni element in SFNM-oxi is +2, while the peaks of SFNM+0.0Fe-red and SFNM+1.2Fe-red overlap with those of Ni foil and NiO. The Ni peaks of both reduced samples are similar, which is in consistence with the XPS and XANES data. It demonstrates that the amount of the Ni exsolution has almost reached the peak value via conventional exsolution due to its lower co-segregation energy. Therefore, combining the XANES, EXAFS results with the XPS data, the Fe and Ni exsolution of both samples can be well discussed. Figure 17). (a) Fe K-edge and (b) Ni K-edge X-ray absorption near-edge structure (XANES) spectra of SFNM-oxi, SFNM+0.0Fe-red and SFNM+1.2Fe-red, together with reference samples Fe-foil, FeO, Fe2O3, Ni-foil, and NiO. Fourier-transformed (c) Fe K-edge and (d) Ni K-edge extended X-ray absorption fine structure (EXAFS) spectra of SFNM-oxi, SFNM+0.0Fe-red and SFNM+1.2Fe-red, together with reference samples Fe-foil, FeO, Fe2O3, Ni-foil, and NiO.

Figure R3 (also Supplementary
We have revised the manuscript accordingly, as shown on L6-8, P17 of the revised manuscript and P23-24 of the revised supplementary information (SI).
Comment #10. In Supplementary Fig. 5, the exsolved particle diameter (or particle size) becomes larger with increasing 'x' value in SFNM + 'x' Fe-red. This would also affect the electrochemical performances of SOFC and/or SOEC.

Authors' response to Comment #10:
We agree with the reviewer's observation. As shown in Supplementary Figs. 5a-d, the TIE-assisted procedure yields an average particle size of around 22 nm within a distribution range of 10-40 nm, which are evidently larger than the corresponding values of 15 nm and 8-22 nm in the SFNM+0.0Fe-red. The increase in the size of exsolved nanoparticles indeed has an impact on the catalytic performance, however, the population of exsolved nanoparticles also increases significantly by ion exchange (Supplementary Fig. 6), thereby expanding the surfaceactive area. From the current density-voltage profiles in Fig. 3a, the overall electrochemical performance of SFNM+1.2Fe-red is superior to that of SFNM+0.0Fe-red. It is in consistence with the results reported in the previous literatures on improving catalytic activity by ion exchange 23 , 24 .
We believe that the changes in both the size and population of exsolved nanoparticles due to the ion-exchangepromoted exsolution contribute to the improved electrochemical performances of SOEC. The effects of the size and population of exsolved nanoparticles from SFNM+ Fe-red ( =0.5, 0.8, 1.2) on the electrochemical performance have been discussed on P8 of the revised SI. Comment #11. In the TGA part (Supplementary Figure 14), the authors stated as "the gradual weight loss below 400 o C can be ascribed to … and decomposition of nitrate". Can the nitrates still exist after synthesis via sol-gel method plus reduction? Moreover, even though the authors wrote Kroger-Vink notation to explain the slope variations with increasing temperature, yet the clear reasons on each slope variations are not clearly elucidated in my position.

Authors' response to Comment #11:
As discussed on the P18-19 of the revised SI, the weight losses of all samples can be ascribed to the detachment of the surface adsorbates and loss of lattice oxygen of P-eNs. In particular, the lattice oxygen loss (formation of oxygen vacancy) above 400 o C is balanced by the valence state change of the B-site transition cations of perovskites 25 (Tables  R1 and R2). According to the valence state changes of Fe, Ni and Mo, the reason for the loss of lattice oxygen at each stage for both samples can be clearly expressed by the Kroger-Vink notation ( is lattice oxygen, •• is oxygen vacancy,   Figure R9. XPS results of (a) Fe (b) Ni (c) Mo for SFNM+1.2Fe-red at 400, 450, 600, 700 and 800 o C. In addition, we thank the reviewer for pointing out the inappropriate term "decomposition of nitrate". We have deleted the phrase of "and decomposition of nitrate" in the revised SI.
Comment #12. How did the authors confirm the exsolved nanoparticle as having Fe:Ni = 1:3 except for the XRD patterns? More clear evidence is required to confirm the ratio of Fe/Ni in Fe-Ni alloy.

Authors' response to Comment #12:
In addition to the XRD results from this work, we have also presented the TEM (with EDS) results of the randomly selected nanoparticles on SFNM+0.0Fe-red ( Supplementary Fig. 4a-c); the calculated average atom ratio of Fe/Ni in the alloy nanoparticles is 0.37 (Supplementary Table 1), which is very close to 1:3 (0.33).
The SFNM-similar perovskites have also been reported in the previous literatures, the exsolved FeNi3 alloy nanoparticles have been reported after a similar reducing treatment in the Reference 26 titled "In situ exsolved FeNi3 nanoparticles on Ni doped Sr2Fe1.5Mo0.5O6−δ perovskite for efficient electrochemical CO2 reduction reaction" and Reference 22.
Comment #13. At the DFT computational details in the methods part, the Gaussian smearing factor is 0.2 eV, which is much high compared to other DFT-conducted literatures. Also, what is the related orbital for Ni and Fe for U calculation? The U values for Ni and Fe seems to be different with other reported literatures.

Authors' response to Comment #13:
Our thanks to the reviewer for raising this question. The sigma values adopted by most literatures for the calculations of metal oxides/perovskites are 0.05, 0.1 and 0.2 27 , 28, 29 , 30 . , We performed three tests with the sigma values being 0.05, 0.1, and 0.2, to firstly optimize the geometric structure of Sr4Fe2Mo2O12 (SFMO), and then to perform the self-consistent calculations (calculation of the total energy), and the results are shown in the Table R3.
For all three sigma values tested (0.05, 0.01, and 0.2), the calculated total energy of the SFMO bulk is very similar (around -274.94 eV), however, the computation time increases as the sigma value decreases. Of note, for sigma = 0.05, the computation time increased 2.4-fold compared to the that of the sigma = 0.2. Based on the results, we chose sigma = 0.2 to obtain reliable value with the reasonable amount of time used. It is also worth mentioning that, in this study, our goal is to calculate the "energy difference" (i.e., exsolution energy, formation of oxygen vacancy) in order to help us to understand the entire reaction process. Therefore, choosing the sigma value of 0.2 and keeping the value consistent throughout the calculations are a reasonable and safe choice for the calculations performed in this study. For the U value, we considered the d-orbitals of Fe and Ni, and we chose the U values based on multiple previous literature results: 31,32 , 33 For example: Fe in LaFeO3 (UFe = 5.1 eV) 34 , Fe in LixFeSiO4 (UFe = 5 eV) 31 , Fe in Sr3Fe1.8Co0.2O7-δ (UFe= 5.3) 35 , etc. In another study, Bouhafs et al. tested the U value from 2 to 8 eV in PrFeO3, which is also a perovskite, and they showed that the calculational results match well with the experimental data (lattice parameters, magnetic moment, band gap, etc.) when using UFe = 5 eV, close to the value of 5.3 that we adopted in this study 33 .
For Ni atom, the literatures with Ni in LaNiO3 (UNi = 6.4 eV) 32 , Ni in LixNiSiO4 (UNi = 6 eV) 31 , doped NiO (UNi = 6.4 eV) 36 all showed a similar U value of Ni to the one adopted in this study.
Therefore, we adopted U=5.3 eV and U=6.2 eV for Fe and Ni, respectively, and believe that this is suitable for the calculations performed in this study.

Note 2: The Materials
Project is an open-access database offering material properties with the structures of more than 35,000 molecules and over 130,000 inorganic compounds. Details for DFT calculations, including the calculations input files for the geometrical optimizations and static calculations, are also provided.
Comment #14. I think that the equation for the calculation of CO production and FECO did not precisely consider the GC measurements.